Monolithic, fully dense silicon carbide material, method of manufacturing and end uses

ABSTRACT

A new silicon carbide material is made following a procedure including hot pressing to provide a finished product having a microstructure with an optimal grain size of less than 7 micrometers. The material exhibits a dominant failure mode of intergranular fracture requiring significant energy for crack propagation. The method of manufacturing is cost-effective by allowing the use of &#34;dirty&#34; raw materials since the process causes impurities to segregate at multi-grain boundary junctions to form isolated pockets of impurities which do not affect the structural integrity of the material. End uses include use as protective projectile-resistant armor.

BACKGROUND OF THE INVENTION

The present invention relates to a monolithic, fully dense, siliconcarbide material, methods of manufacturing the material and end uses.The main use for the present invention consists of use asprojectile-resistant armor for mounting on vehicles or other objects anddevices which must be resistant to penetration by various projectilessuch as, for example, kinetic energy or shaped chargedpenetrators/projectiles and ballistic fragments.

The following prior art is known to Applicant:

U.S. Pat. No. 3,592,942 to Hauck et al. discloses the use of siliconcarbide as an armor material. The present invention differs from theteachings of Hauck et al. as contemplating a particular grain size andstructural mode of crack propagation and resistance thereto nowheretaught or suggested by Hauck et al.

U.S. Pat. No. 3,765,300 and 3,796,564 to Taylor et al. teach a densecarbide composite for armor and abrasives. Taylor et al. disclose that". . . it is essential that the granular boron carbide in the initialmixture have a maximum particle size of about 300 microns or less,although coarser material may be employed to make composite bodiesuseful for less demanding purposes." Taylor et al. further disclose that"The modulus of rupture may be as low as 10,000 psi (700 kg/cm²),especially where granular boron carbide with a particle size greaterthan about 300 microns is included in the mix . . .". As such, Taylor etal. fail to contemplate the grain size disclosed herein nor the methodof manufacturing nor the intergranular fracture and resistance theretodisclosed herein.

U.S. Pat. No. 3,793,648 to Dorre et al. discloses a bullet resistingarmor including a ceramic layer. This patent is believed to be of onlygeneral interest concerning the teachings of the present invention.

U.S. Pat. No. 3,859,892 to Coes discloses a composite ceramic armorincluding disclosure of the use of ceramic plates made of materials suchas "polycrystalline alumina, silicon carbide or boron carbide". Thispatent is believed to be of only general interest concerning theteachings of the present invention.

U.S. Pat. No. 3,977,294 to Jahn discloses composite armor and method.This patent contemplates a composite laminate material including layersof graphite and ceramic materials adhered together through the use of anadhesive. The ceramic materials disclosed herein are nowhere taught orsuggested by Jahn.

U.S. Pat. No. 4,692,418 to Boecker et al. discloses a sintered siliconcarbide/carbon composite ceramic body having fine microstructure. Thispatent discloses making of a silicon carbide material and, thereafter,infiltrating a source of carbon therein to provide the composite body.The present invention differs from the teachings of Boecker et al. ascontemplating a monolithic silicon carbide material possessing uniqueproperties and made in a manner not contemplated by Boecker et al.

U.S. Pat. No. 4,693,988 to Boecker et al. discloses a single phasesilicon carbide refractory. The Boecker et al. material consists of apressureless sintered product made from starting materials wherein acoarse fraction as well as a fine fraction are present. Boecker et al.disclose that the coarse fraction has a particle size ranging between210 and 340 micrometers, huge particles as compared to thosecontemplated herein. In Boecker et al., all particles are carriedthrough to the final microstructure, that is, they do not reduce in sizeduring the sintering process. Ballistic test results show that suchgrains incorporated within a SiC microstructure would clearly exhibittransgranular fracture when impacted by a projectile as compared to thepresent invention wherein intergranular fracture would occur butmechanisms to resist such fracture exist. The inventive grain size(material microstructure) disclosed in this patent application ispreferably equal to or less than 7 micrometers which is required tofacilitate intergranular fracture. That is, experimentation has revealedthat silicon carbide grains larger than 7 micrometers generally exhibittransgranular fracture and grains smaller than 7 micrometers generallyexhibit intergranular fracture. As such, the present invention clearlydiffers from the teachings of Boecker et al.

U.S. Pat. No. 4,813,334 to Bloks et al. discloses armor plates includinglayers of ceramic material and fiber reinforced plastic separated bymetallic plates. This patent is believed to be of only general interestconcerning the teachings of the present invention.

U.S. Pat. No. 4,876,941 to Barnes et al. discloses a composite forprotection against armor-piercing projectiles. The disclosed material isa ceramic composite comprising titanium boride combined with aluminumnitride. Hot pressing techniques are employed in the manufacture of thismaterial. The present invention differs from the teachings of Barnes etal. as contemplating a monolithic silicon carbide evidencing a mode offailure defined as intergranular in nature. While aluminum nitride isemployed in the process of manufacturing the inventive ceramic, it isonly used as a densification aid and in proportion much smaller than theproportion of silicon carbide which is employed. In a composite, such asthat which is disclosed by Barnes et al., the aluminum nitride is anintegral part of the microstructure, that is, aluminum nitride grainsare present and can be specifically identified as aluminum nitride. Incontrast to this, in the present invention, aluminum nitride behaves asa densification aid. After processing, one may not identify specificaluminum nitride grains. Aluminum nitride is not an integral part of themicrostructure and, chemically, the finished ceramic body does not showthe presence of aluminum nitride.

K. Nakamura and K. Maeda, in Silicon Carbide Ceramics, Volume 2, Editedby S. Somiya and Y. Inomata, Elsevier Applied Science, 1991, disclosehot-pressed SiC ceramics. These investigators have demonstrated ahot-pressed silicon carbide material using aluminum nitride (AlN) as aprocessing aid with a Weibull modulus of 13.8. The inventive siliconcarbide material disclosed herein contemplates a Weibull modulus withina range of 18 to 30, much higher than Nakamura et al. which makes itpossible to produce components with outstanding performancecharacteristics and exceptional reliability where prior state-of-the-artmaterials could not be used. The present invention also differs from theteachings of Nakamura et al. as contemplating densification techniquesand relation between weight of AlN addition to SiC powder surface areanowhere taught or suggested therein.

SUMMARY OF THE INVENTION

The present invention relates to a fully dense, monolithic form ofsilicon carbide material, method of manufacturing and end uses. Thepresent invention includes the following interrelated objects, aspectsand features:

(A) The particular silicon carbide ceramic material which is employed ismade up of constituent substances including at least 92%, by weight,silicon carbide powder of the alpha or beta type. Powder specifications,in the preferred embodiment, also include from 0.01% up to 2% iron, from0.10% up to 2% free carbon, from 0.01% up to 1.5% aluminum, from 0.02%up to 3% silicon dioxide, from 0.01% to 2.5% oxygen and from 0.01% up to0.15% free silicon, with these figures being by weight. Other cationimpurities such as Ca, Mg, Ti, Na, W, etc. are also permissible as wellas anion impurities such as N. These impurities are present in mostcommercially available powders on a trace basis, e.g., 0.01% or less.

(B) Concerning particles making up the powder, average particle sizeshould not exceed 3 micrometers, with maximum particle size being lessthan 10 micrometers and with the optimal maximum particle size being nogreater than 7 micrometers. Furthermore, the powder surface area shouldexceed 5 m² /g. As will be described in greater detail hereinafter, dueto the unique process of manufacture of the inventive material,relatively impure silicon carbide powder may be employed.

(C) In the manufacture of the inventive material, as will be describedin greater detail hereinafter, a dry procedure or, alternatively, a wetprocedure may be employed. After one or the other of these procedureshas been employed, hot pressing of the processed powder is accomplishedunder precise conditions of temperature and pressure through a regimenfollowed by a cooling down procedure which results in manufacture of thefinished inventive material.

(D) The inventive material has been tested in terms of its resistance topenetration by projectiles. Such testing has revealed great resistanceto penetration, to an extent far exceeding performance of prior artsilicon carbide materials. In fact, as will be described in greaterdetail hereinafter, the inventive material shows such a degree ofresistance to penetration that the thickness of ballistic tile can bereduced dramatically and thereby reducing the weight of the associatedarmor system. Prior art projectile resistant silicon carbide materialsrequire thick panels since their ballistic performance is sub-standardto that provided by this inventive material. The fact that the inventivematerial may be employed with a thin metallic backing plate allows theend user to save a great deal of weight since, for example, RHA steelweighs 41.6 pounds per square foot in a one inch thickness. This steelis a preferred backing material in prior art armor systems. Theinventive ceramic material may be used in thinner layers than prior artceramic armor materials and backing plate materials may be ferrous ornon-ferrous (aluminum), glass, glass composites, KEVLAR or a combinationthereof.

As such, it is a first object of the present invention to provide amonolithic silicon carbide material, method of manufacturing and enduses.

It is a further object of the present invention to provide such amaterial having an optimal grain size of less than 7 micrometers.

It is a still further object of the present invention to provide a fullydense (98.5% of theoretical density or greater), high strength, highlyreliable material characterized by a Weibull modulus of greater than 18.

It is a still further object of the present invention to provide such amaterial made using hot pressing techniques at relatively lowtemperatures.

It is a still further object of the present invention to provide such amaterial which fractures in an intergranular fashion and resists suchfractures.

It is a still further object of the present invention to provide such amaterial that contains segregated grain boundaries.

It is a yet further object of the present invention to provide such amaterial which resists penetration of projectiles to an extent obviatingthe use of "thick" ceramic panels.

It is a still further object of the present invention to provide such amaterial that exhibits a high fracture toughness, an exceptionally highdegree of thermal conductivity and a high electrical resistivity.

It is a still further object of the present invention to provide such amaterial which may be used as a projectile resistant armor as well as inother applications such as in the environment of optical and mediastorage discs materials.

These and other objects, aspects and features of the present inventionwill be better understood from the following detailed description of thepreferred embodiment when read in conjunction with the appended drawingfigures.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 shows a photomicrograph of a fracture surface of a commerciallyavailable monolithic silicon carbide material as known in the prior art,e.g., hot pressed using boron and carbon as sintering aids.

FIG. 2 shows a photomicrograph of the fracture surface of the siliconcarbide material made in accordance with the teachings of the presentinvention.

FIG. 3 shows a graph of depth of penetration versus velocity ofprojectile for monolithic rolled homogenous annealed steel, which iscommonly used as a backing material in projectile resistant armor.

FIG. 4 shows a graph of depth of penetration for various armor materialsincluding the material disclosed herein.

FIGS. 5 and 6 comprise photomicrographs showing the fracture surface ofa monolithic, fully dense silicon carbide material made in accordancewith the teachings of Example 1.

FIG. 7 comprises a photomicrograph showing the intergranular fracturesurfaces of a ceramic material made as set forth in Example 2.

FIG. 8 comprises a photomicrograph showing the fracture surfaces ofceramic material made in accordance with the parameters of Example 3.

FIG. 9 is a graph showing Weibull modulus variation with variation ofAluminum Nitride (sintering aid) content.

SPECIFIC DESCRIPTION OF THE PREFERRED EMBODIMENT

Reference is first made to FIG. 1 which depicts a photomicrographshowing the fracture surface of a prior art silicon carbide material.This photomicrograph is presented in this application as an example ofthe prior art and to place the present invention in its properperspective.

As should be understood, when reviewing FIG. 1, the fracture surface ofthe prior art silicon carbide depicted exhibits transgranular fracture.What is meant by this is that when the material depicted in FIG. 1fractures, it "cleaves" linearly. Thus, when a projectile strikes thesilicon carbide material shown in FIG. 1, limited resistance topenetration of the projectile is exhibited due to the linear nature ofthe fracture.

In contrast to the teachings of the prior art silicon carbide materialas illustrated in FIG. 1, FIG. 2 depicts a photomicrograph of thefracture surface of a monolithic, fully dense silicon carbide materialmade in accordance with the teachings of the present invention. Asshould be understood from FIG. 2, the fracture surface of the inventivesilicon carbide material exhibits intergranular fracture. What thismeans is that when the silicon carbide material shown in FIG. 2fractures, it does so non-linearly in an undulating path back and forthbetween the various grains contained therein. The non-linear nature offractures which are formed in the inventive material as shown in FIG. 2renders the material quite effective in resisting penetration byprojectiles Forces imposed upon the material through impact from aprojectile are spread laterally with respect to the direction of travelof the projectile with projectile forces being spread in such a mannerthat energy is rapidly dissipated in directions other than the directionof travel of the projectile. This is to be compared, again, with resultsobtained through the use of the silicon carbide material illustrated inFIG. 1 as an armor material wherein impact by a projectile causes thematerial to cleave and fail catastrophically, thereby allowing theprojectile to relatively more easily travel therethrough.

In armor, it is common to provide an armor material backed by a backingsuch as rolled homogenous annealed steel. FIG. 3 depicts a graph ofdepth of penetration into rolled homogenous annealed steel byprojectiles versus the velocity in feet per second of the particularprojectile. As should be understood from FIG. 3, the depth ofpenetration into the steel base is approximately proportional to theincrease in velocity of the projectile.

FIG. 4 shows a graph of depth of penetration for various armor materialsbacked by ballistic grade aluminum. Ballistic grade aluminum weighs14.04 pounds per square foot at a one inch thickness. The depth ofpenetration depicted in FIG. 4 above zero is the depth of penetrationinto the aluminum. Thus, any materials which show any depth ofpenetration in the graph depicted in FIG. 4 have completely penetratedthrough the identified armor material. The projectile which was used intesting each material is a Russian 14.5 mm BS 41 tungsten carbide coredprojectile fired at a muzzle velocity of about 3200 to 3400 feet persecond.

As shown in FIG. 4, use of armor material made in accordance with theteachings of the present invention resulted in zero penetration into thebacking aluminum. The other eight materials, all of which comprise someform of silicon carbide made in processes differing from the processemployed herein exhibited some degree of penetration into the backingaluminum, from 0.35 inches in the case of material A to as much as 1inch for the material identified by the letter H.

Rolled homogenous annealed steel, a typical backing material, weighs41.6 pounds per square foot at a one inch thickness, almost three timesthe weight for ballistic grade aluminum of equal dimensions. It shouldbe understood that use of the present invention as an armor materialwill save a significant amount of weight over prior art armor materials.Thus, for example, were armor material to be made of the hot pressedsilicon carbide material identified by the letter A in FIG. 4, one wouldneed to employ a rolled homogenous annealed steel base in thickness andpenetration resistance equal to at least 0.35 inches of ballistic gradealuminum performance. Based upon the weight figures noted above, use ofthe present invention would allow use of a much thinner base material.Thus, it should be clear that large weight savings would accrue whenutilizing the teachings of the present invention to armor, for example,a tank, personnel carrier or ship. Also given the penetration resistanceof the present invention, it may be applied with less thickness thanother ceramic armor materials, thereby saving even more weight.

In an important aspect of the present invention, aluminum nitride isutilized as a densification agent. As explained above in the BACKGROUNDOF THE INVENTION, it is known to utilize aluminum nitride in combinationwith silicon carbide to form a composite. However, Applicant's use ofaluminum nitride in the inventive process does not result in theformation of a composite, for many reasons including the smallproportion of the aluminum nitride as compared to the amount of siliconcarbide employed. Additionally, the Nakamura et al. publicationdiscloses use of aluminum nitride as a sintering aid at concentrationsof 0.5% to 4%, by weight, of silicon carbide. However, Nakamura et al.fail to contemplate the significance of the ratio between the weight ofaddition of aluminum nitride and the total surface area of the siliconcarbide powder, nor do Nakamura et al. contemplate the densificationtechniques disclosed herein nor the resulting enhanced Weibull modulusnor the ability to form segregated grain boundaries.

Aluminum nitride (AlN) makes the surfaces of the silicon carbide"active" and increases diffusional kinetics via the grain boundariesthereof. Through experimentation, Applicant has discovered that theoptimum amount of aluminum nitride to be added to the silicon carbide isbased upon the surface area of the silicon carbide powder. In thisregard, the total amount of surface area of the silicon carbide powdergenerally defines the grain boundary surface area and, therefore, theamount of area which must be "activated" to optimize the finishedceramic product. Through experimentation, Applicant has found that theoptimum content of aluminum nitride is 0.10-0.25%, by weight, aluminumnitride for each m² /g of surface area of silicon carbide.

Applicant has discovered that below a minimum weight addition level ofaluminum nitride with respect to the surface area of the silicon carbidepowder, there are insufficient quantities of aluminum nitride availableto activate surface and/or bulk diffusional kinetics to efficiently andtotally densify the silicon carbide material. That is, the sinteringevent is inefficient and will produce ceramic bodies that do not meetdensity requirements of the finished products. In the case of the armorapplications as disclosed herein, the finished ceramic body must be asclose to theoretical density as possible, certainly at least 98.5% oftheoretical density. Ceramic bodies that contain a relatively highdegree of porosity (low density) are characterized by low Weibullmodulii because pores and pore clusters comprise the dominant flawpopulation.

For example, the use of 0.05% by weight aluminum nitride for each m² /gof surface area of silicon carbide powder where the silicon carbidepowder is of the alpha type and has 12 m² /g surface area will result,in accordance with the hot pressing conditions disclosed hereinafter, ina finished ceramic body having a density of about 87% of theoreticaldensity. Such a ceramic body is characterized as having a relatively lowMOR and is represented by a low Weibull modulus, less than 10. Fracturesurface examination of such a body shows that cracks initiate at largepores and at regions of pore clustering. Above the maximum permissibleratio of addition by weight of aluminum nitride to surface area ofsilicon carbide powder, to-wit, above 0.25% by weight aluminum nitridefor each m² /g of surface area of silicon carbide, a surplus of aluminumnitride exists in unreacted form and remains at the grain boundaries orforms compounds based upon aluminum nitride which constitute a secondaryintergranular phase. This phase may be crystalline, amorphous or acombination thereof. Intergranular phases and/or large particles ofunreacted aluminum nitride form defects which are deleterious to theperformance of the material and produce materials, again, with lowWeibull modulii.

The aluminum nitride activates the silicon carbide grain surfaces in asolubility reaction. At the hot pressing temperatures disclosedhereinafter, the solubility of aluminum nitride is confined to thesurfaces of the silicon carbide particles. Furthermore, the solubilityof aluminum nitride in silicon carbide is also limited. Therefore, whenthe solubility limit is exceeded, that is, additional aluminum nitridecannot be absorbed into the silicon carbide structure, the superfluousaluminum nitride remains in the grain boundaries in an unreacted form orreacts with other grain boundary materials, for example, silica, to formintergranular materials.

Applicant has selected aluminum nitride for use as a densification aiddue to its efficiency as a sintering aid and due to reduced thermalenergy required for densification. That is, AlN enables SiC mixtures tobe hot pressed at reduced temperatures thereby decreasing exaggeratedgrain growth during hot pressing, enabling better control over grainsize in the final developed microstructure, e.g., grains of 7micrometers or less. Reduced temperatures deter grain growth duringprocessing. Most standard grades of aluminum nitride powder have anaverage particle size of 4.0 micrometers or less and, as such, areacceptable as a sintering aid. It must be noted, however, that aluminumnitride is hydroscopic and, as such, requires handling in a manneravoiding contact with water or water vapor. In light of this factor, theinventive process may not involve the use of water or water vapor unlessthe aluminum nitride powder is coated with a protective film. However,if the aluminum nitride powder is coated with a protective film topreclude hydrolysis, then water may be used as a milling/homogenizingmedia.

FIG. 9 shows the significantly enhanced Weibull modulus which results inaccordance with the teachings of the present invention, particularlymaintaining the proportion of AlN (sintering aid) within the disclosedrange. FIG. 9 shows a data point representative of the Nakamura et al.material.

Through experimentation, Applicant has found that impurities in the rawmaterials do not harm the eventual performance of the finished siliconcarbide material. This is an important factor since highly pure rawmaterials are significantly more expensive than relatively "dirty"grades of raw materials. The inventive process which is employed inmaking the silicon carbide material disclosed herein causes impuritiesto segregate at grain boundaries and therefore these impurities do nothave a direct impact on the intrinsic strength of the present invention.The grain boundaries segregate the cation impurities, formingdiscontinuous pools of impurities, generally submicrometer in size,which are inert within the finished product material. The inventivematerial has been made using silicon carbide powder at prices less than$1.50 per pound, much less than the price per pound of more highly puresilicon carbide powder.

In a further aspect, one may apply Weibull analysis to describe thestrength variation of various ceramic armor materials. Weibull analysisis based upon the "weakest link of the chain" theory whereby thestrength of a ceramic material strongly depends upon the size and shapeas well as frequency of internal defects. These internal defects willreside in areas where failures are more likely to occur. In performingWeibull analysis, sufficient statistically significant amounts of testspecimens must be prepared and tested. Uniform standards for dimensionsare complied with and flexural strength is calculated through the use ofa loading and testing fixture in a manner known to those skilled in theart. A plurality of specimens are tested, one-by-one, and may be rankedin ascending order of strength with each specimen being assigned aprobability of failure according to its corresponding rank in order.

The probability of failure of any one specimen is calculated by usingthe equation: ##EQU1## where F is

The probability of failure, N is the total number of specimen is and nis the rank of the specific specimen being tested.

Weibull analysis shows the relationship between the probability offailure and the strength of a specimen through the use of the followingequation: ##EQU2## where ln is the natural logarithm, F is theprobability of failure, S is the strength of the specific bar, So is thecharacteristic strength and m is the Weibull modulus.

In general, a high Weibull modulus means narrow strength distribution.The characteristic strength represents the strength having 63.2% offailure probability. Thus, the Weibull modulus is calculated bylinear-least-square fitting of the data points calculated through theuse of the latter-mentioned formula for the relationship between theprobability of failure and strength of each specimen.

Table A displays the various properties of five conventional siliconcarbide materials as well as the inventive silicon carbide materialincluding display of the Weibull modulus. As is seen in Table A, theWeibull modulus for the present invention is greater than 20 whereas thenext best material exhibits a Weibull modulus of no greater than 11.This fact combined with the fact that the measured flexural strength forthe inventive silicon carbide material generally exceeds 100,000 psi,should make clear the significant improvement in performance of thepresent invention as compared to the prior art. While the Nakamura etal. material also exhibits a flexural strength in excess of 100,000 psi,Nakamura et al. require the use of extremely expensive, highly pureconstituent ingredients to achieve this level of flexural strength. Bycontrast, as disclosed herein, the present invention maintains allinventive parameters even where relatively "dirty" constituent materialsare employed.

                  TABLE A                                                         ______________________________________                                        SiC type α-SiC (SC-501)*                                                                      α-SiC (SASC)**                                                                      β-SiC***                               ______________________________________                                        Designation                                                                            Hot Pressed  Sintered    Sintered                                    Density  3.2          3.11-3.13   3.12-3.15                                   (g/cm.sup.3)                                                                  Free C   0.24         1.09        0.89                                        (wt %)                                                                        Oxygen   N/A          0.059       0.044                                       (wt %)                                                                        MOR test RT****       RT          RT                                          (°C.)                                                                           1200         1370        1370                                        Mean MOR 750          380 ± 46 423 ± 55                                 (Mpa)    750          307 ± 50 388 ± 59                                 Weibull  13.5         9.6         9.5                                         (m)      13.5         7           8.1                                         ______________________________________                                                              αSiC (Type                                                                          αSiC (Type                            SiC type SiC/Si (KXO1)**                                                                            A)*****     B)*****                                     ______________________________________                                        Designation                                                                            Reaction-bonded                                                                            Hot Pressed Hot Pressed                                 Density  2.89-2.91    3.16-3.22   3.20-3.24                                   (g/cm.sup.3)                                                                  Free C   0.22         1.20        2.0% Max                                    (wt %)                                                                        Oxygen                                                                        (wt %)   0.11         0.10        0.64                                        MOR test RT           RT          RT                                          (°C.)                                                                           1200         1200        1200                                        Mean MOR 385 ± 58  520 ± 60 670 ± 42                                 (Mpa)    415 ± 51  510 ± 54 648 ± 49                                 Weibull  8            8           26                                          (m)      9.2          9.5         24                                          ______________________________________                                         *Hitachi (Nakamura et al. material) **Carborundum Co. ***General Electric     ****Room Temperature *****Cercom Inc.                                    

In conducting the process to produce the inventive alpha SiC materials,blends of commercial SiC powders and AlN powders are prepared,homogenized, and hot-pressed in accordance with the parameters describedbelow.

A fine-grained silicon carbide powder as defined below by particle sizeis used. Generally, the average particle size should not be greater than3 micrometers to result in the new alpha SiC composition of the currentinvention which predominantly exhibits intergranular fracture. A powderpurity specification is prepared in accordance with the parametersdisplayed below in Table B as follows:

                  TABLE B                                                         ______________________________________                                        Constituent Material                                                                         % By Weight in Mixture                                         ______________________________________                                        Si C           92% minimum                                                    Fe             0.01-2.0                                                       Free C         0.10-2.0                                                       Al             0.01-2.0                                                       SiO2           0.02-3.0                                                       Oxygen         0.01-2.5                                                       Free Si         0.01-0.15                                                     Other, Individually                                                                          Trace < .01                                                    ______________________________________                                    

The powder purity specification shown in Table B is also required, inaccordance with the teachings of the present invention, to comply withadditional requirements. The particles included in the mix may notexceed 10 micrometers in diameter with an average particle diameter of0.3 to 3 micrometers being optimal. The cumulative surface area of theparticles should be within the range of 3 to 20 m² /g. This factor maybe measured in a manner known to those skilled in the art.

With the silicon carbide powder purity selected in accordance with TableB and the above additional specifications including an appropriateselection of AlN powder, powder formulation and processing of powdermixtures may proceed, in accordance with the teachings of the presentinvention, in a powder homogenizing and comminution procedure which mayeither comprise a dry procedure or a wet mill procedure.

In the dry procedure, a master mix is prepared with the weightsindicated below being exemplary. The master mix techniques set forthbelow are used to enable introducing of small amounts of secondaryconstituent powders (on a weight % basis) into a main powder constituentin a homogeneous basis particularly where the powder process is a dryprocess. These techniques are generally known in the art.

In a 25 liter polypropylene bottle, the following constituent substancesare added: 15 kg of a grinding media consisting of a fully denseporcelain body made up of about 90% Al₂ O₃ and the balance SiO₂, i.e., a1/2 inch diameter rod, 5100 g of alpha-type SiC powder and 900 g ofaluminum nitride powder.

When this mixture has been prepared, homogenization is accomplished byadding 6500 milliliters of methanol, ethanol or isopropyl alcohol (allanhydrous), milling for approximately 5-8 hours, subsequently pan dryingthe mixture for 20-24 hours at about 90 degrees C. and, thereafter, drymilling for one half hour with the 1/2 inch diameter grinding rod.

Thereafter, the mixture is blended in the proportion of 81 Kg of siliconcarbide combined with 9 kg of the master mix as described above untilhomogeneous.

As described above, an alternative to the dry procedure comprises a wetmill process. In this process, in a 25 liter polypropylene bottle, 15 kgof grinding media, as described above, is added and a powder charge isprovided including about 6649 g of silicon carbide and 101 g of aluminumnitride. 7500 milliliters of alcohol (anhydrous) is added and wetmilling is carried out to achieve homogeneity between the constituentingredients followed by drying to remove the carrier (alcohol), followedby dry milling to break up the agglomerates and screening through 30mesh.

After the powder has been processed in accordance with either one of theprocedures described above, the intermediate product is then ready forintroduction into a hot press mold cavity and subsequent hot pressing tofull density. To commence the hot pressing procedure, the hot pressingchamber is evacuated to 1.5 TORR and then the chamber is either backfilled with nitrogen or argon or the entire procedure is run undervacuum. Table C, displayed below, sets forth the preferred process stepsfor the hot pressing multiple SiC plates dimension 12"×12"×2". Theschedule shown is representative of a large commercial induction-heatedvacuum hot press and has been adapted accordingly. The process stepstake into consideration such concerns as thermal mass, achieving thermalequilibrium, outgassing of raw materials, etc.

                  TABLE C                                                         ______________________________________                                                   Pressure    Parameters (Ramp rate,                                 Temperature                                                                              Part        powder setting, hold                                   (Degree-C.)                                                                              (psi)       time, and atmosphere                                   ______________________________________                                        Room Temp   500                                                               1630        500        Hold 2 hours, then                                                            increase @ 1° C./minute                         1820       1000        Hold 40 minutes, then                                                         increase @ 1° C./minute                         1895       1500        Hold 40 minutes, then                                                         increase @ 1° C./minute                         1900       2000        Hold 40 minutes                                        *1900      2500        Hold 1.5 hours, then                                                          increase @ 1° C./minute                         2040-45    2500        Hold until ram movement                                                       essentially stops, not                                                        to exceed 5 hours                                      ______________________________________                                         *80% of drop must come out of system at this temperature, or full density     will not be achieved. That is, a closed porosity situation must be            established, or volatization of metals/metaloids occurs as temperature        increases  these materials assist in densification. This is the essence o     the hot pressing procedure. That is, the hot pressing procedure must be       adapted so that at least 80% density (of theoretical) is obtained at          temperatures equal to or less than 1900° C. Temperatures above         1900° C. are used to achieve a fully dense state. That is, to          remove the remainder of void volume. Temperatures above 2050° C.       are not used so as to preclude exaggerated grain growth.                 

Through experimentation, Applicant has found that at approximately 80%of theoretical density, most porosity within the ceramic body is closed.Furthermore, above 1900° C., sintering aids and cation impurities beginto volatize vis.a.vis decomposition. In order to effectively densify theceramic body to near theoretical density, sintering aids must remainwithin the body. In a closed porosity condition, sintering aids cannotleave the body by volatization, that is, the sintering aids are capturedwithin the body. Therefore, in accordance with the teachings of thepresent invention, it is important that a closed porosity condition beachieved before the critical temperature of 1900° C. is exceeded.

After the procedure set forth in Table C has been followed, theresulting product, if of large cross-sectional area, must be cooled downin a controlled fashion or it will be subject to thermal shock. The cooldown procedure is determined emperically and is adapted to the equipmentused and the part(s) being manufactured. Generally, the hot pressingpressure is allowed to decay linearly with decreasing temperature withall pressure being removed at 1500° C. Temperature decay rate is themost important consideration wherein temperature is reduced at a ratecompatible with minimal thermal stress developing in the cooling SiCproduct. Where small cross-section articles are involved, they may becooled by allowing the furnace in which they are contained to naturallycool. The finished product described herein will result from followingof these procedures.

The following are examples of practicing of the inventive process.

EXAMPLE 1

Commercially available Ultrafine Alpha-Type Silicon Carbide Powder wasused as the starting powder. This grade has an average particle size of0.7 micrometers and the measured surface area of the particles was 13m²/g. Given this high surface area to weight ratio, 3% by weight ofaluminum nitride was used as a sintering aid. The powder was processedin anhydrous alcohol to prevent aluminum nitride hydrolysis. Neoprenelined Jar mills with Al₂ O₃ grinding media were used to homogenize thebatched powder. A time increment of 6 hours at a slurry viscosity of 200centipoise produced a well dispersed powder blend as determined byoptical microscopy. All powders were prepared for cold pressing by pandrying, dry milling and subsequent screening in accordance with theteachings of the process as described hereinabove. Preforms wereproduced by pressing in steel dies to a green density of approximately45% of theoretical. A die body and die cavity were prepared using theinventive silicon carbide procedures as described herein. Hot pressingwas conducted at 3500 psi and at a temperature of 2000° C. with 85% ofdensity being developed below or at a temperature of 1900° C.Thereafter, three hours of treatment at the ultimate hot pressingtemperature was required to reach full density. Evaluation of thefinished microstructure occurred with particular emphasis on evaluationof representative fracture surfaces. These were examined by optical andscanning micrography. Etched samples showed no evidence of porosity anda generally equiaxed fine-grained microstructure. Grain size wasmeasured by using the ASTM E-112 standard and was calculated to be 2.16micrometers. FIGS. 5 and 6 are photomicrographs of representativefracture surfaces at 3000× and 5000× magnification, respectively. Theydistinctly show extensive intergranular crack propagation. Intergranularfracture corresponds to a crack deflection failure model and requiressignificant energy for crack propagation, thereby increasing thefracture toughness of the overall material. The physical properties ofthe finished ceramic material were measured or calculated as follows:density was measured by water immersion to be 3.219 g/cm³, slightlyexceeding the calculated theoretical density; the elastic and shearmodulii are 64.22 Mpsi and 28.06 Mpsi, respectively; Poisson's ratio wascalculated to be 0.14; Knoop hardness using 0.5 kg loading wascalculated to be 2747 plus or minus 63 kg/mm₂. No evidence of porositywas found which is supported by the measured density.

EXAMPLE 2

A monolithic, fully dense silicon carbide was prepared usingcommercially available silicon carbide having a surface area to weightratio of 12 m² /g. The powder was processed in conjunction with 1.5%, byweight, aluminum nitride with wet milling in anhydrous isopropyl alcoholfor 8 hours. Hot pressing was conducted at a temperature of 2000° C. anda pressure of 3000 psi for 2 hours. Etched polished samples wereexamined by optical microscopy. The samples showed no evidence ofporosity and a generally equiaxed fine-grained microstructure. Theaverage grain size was about 2 micrometers with very few grains largerthan 7 micrometers. The samples showed predominantly intergranularfracture surfaces as shown in FIG. 7 which comprises a photomicrographat 3000× magnification. The samples were fully dense with a density ofabout 3.2 g/cm³ and the elastic and shear modulii were about 66 Mpsi and29 Mpsi, respectively. Poisson's ratio was calculated to be about 0.15and Knoop hardness using 0.5 kg loading was about 2800 kg/mm². Fracturetoughness was calculated to be 4.5 MPa m^(1/2). The oxygen content ofthe finished product was less than 2% by weight.

EXAMPLE 3

Monolithic, fully dense silicon carbide material was made using, as theraw material, commercially available silicon carbide having a surfacearea to weight ratio of 15 m² /g. 1.5% by weight of aluminum nitride wasmixed with the silicon carbide material and the mixture was processedusing wet milling in a ketone solvent [methyl ethyl ketone (MEK)] for 8hours. The die body and die cavity were prepared using the inventivesilicon carbide procedures. Hot pressing was conducted at a temperatureof 2000° C. and a pressure of 3000 psi for 2 hours. For all samples,diffraction results show no evidence of an amorphous phase or acrystalline phase other than silicon carbide. Etched polished sampleswere examined by optical microscopy with all samples showing no evidenceof porosity and generally equiaxed fine-grained microstructure. Theaverage grain size was calculated to be 1.6 micrometers with very fewgrains larger than 7 micrometers. All samples show predominantlyintergranular fracture surfaces indicating that there was extensivecrack deflection during failure. This type of failure mode generallyleads to high fracture toughness. The density of the finished materialswas about 3.20 g/cm³, as should be expected giving the lack of porositynoted above. The measured elastic and shear modulii were about 59 and 26Mpsi, respectively, with Poisson's ratio being calculated to be about0.15. Knoop hardness using 0.5 kg loading was about 2800 kg/mm² whilefracture toughness was about 3.9 MPa m^(1/2). The lower fracturetoughness (as compared to the material in Example 2) is attributable tomore grains showing transgranular fracture, although the fracturesurface is predominantly intergranular in nature. The oxygen content wasabout 2% by weight. FIG. 8 shows the fracture surface of the ceramicmaterial made in accordance with the parameters of Example 3, showingpredominantly intergranular fracture surfaces. FIG. 8 comprises aphotomicrograph at 3000× magnification.

The materials made in accordance with the teachings of Examples 1, 2 and3 each have a calculated Weibull modulus of greater than 20, more thantwice the corresponding Weibull modulus for other known silicon carbidematerials. As such, the inventive materials are quite suitable for useas projectile resistant armor.

Regardless of whether the initial silicon carbide material is of thealpha type or beta type, the finished product comprises an alpha typesilicon carbide ceramic material. As should be understood, beta typesilicon carbide converts to alpha type silicon carbide at about1850°-1900° C., a temperature exceeded in all applications and examples.

Accordingly, an invention has been disclosed in terms of a material, theprocess of manufacturing the material and its end uses which fulfilleach and every one of the objects of the invention as set forthhereinabove and provide a new and useful invention of great novelty andutility.

Of course, various changes, modifications and alterations in theteachings of the present invention may be contemplated by those skilledin the art without departing from the intended spirit and scope thereof.As such, it is intended that the present invention only be limited bythe terms of the appended claims.

I claim:
 1. A monolithic alpha-type silicon carbide finished material,made from a powder mix consisting of, by weight, at least 92% SiliconCarbide, from 0.01% to 2% Iron, from 0.01% to 1.5% Aluminum, from 0.01%to 2.5% Oxygen and from 0.02% to 3% Silicon Dioxide, said powder mixadditionally including a sintering aid consisting of 0.10% to 0.25%, byweight, Aluminum Nitride for each m² of Silicon Carbide per gram ofSilicon Carbide, said powder mix including particles up to 10micrometers in diameter, said finished material possessing the followingproperties:a) substantially non-porous; b) near theoretical density; c)average grain size about 3 micrometers in diameter with maximum grainsize preferably no more than 10 micrometers in diameter; d) fracturesproceed in an intergranular non-linear manner; e) Weibull modulusgreater than
 18. 2. The material of claim 1, further including from 0.1%to 2.0% free Carbon.
 3. The material of claim 1, further including from0.01% to 1.5% free Silicon.
 4. The material of claim 1, wherein grainsthereof do not exceed 7 micrometers in diameter.
 5. The material ofclaim 1, wherein a ratio of area of particles in square meters to weightin grams falls within a range of 3 to 20 m² /g.
 6. The material of claim1, wherein said near theoretical density is at least 98.5% oftheoretical density.